Borophenes, boron layer allotropes and methods of preparation

ABSTRACT

A method of preparing an atomically-dimensioned elemental boron allotrope includes providing a substrate at a temperature greater than about 200° C.; generating elemental boron vapor from a solid elemental boron source; and contacting said substrate with said boron vapor for at least one of a rate and at a pressure sufficient to deposit on said substrate a boron allotrope comprising an elemental boron layer comprising a boron atomic thickness dimension, said method under negative pressure.

CROSS-REFERENCE TO RELATED PATENT APPLICATIONS

This application is a divisional application of and claims the benefitof U.S. patent application Ser. No. 15/430,885, filed Feb. 13, 2017, nowallowed, which claims priority to and the benefit of U.S. ProvisionalPatent Application No. 62/388,970, filed Feb. 12, 2016, which areincorporated herein in their entireties by reference.

STATEMENT AS TO RIGHTS UNDER FEDERALLY-SPONSORED RESEARCH

This invention was made with government support under DE-AC02-06CH11357and DE-FG02-09ER16109 awarded by the Department of Energy; DMR-1121262and DGE-0824162 awarded by the National Science Foundation; andN00014-14-10669 awarded by the Office of Naval Research. The governmenthas certain rights in the invention.

BACKGROUND OF THE INVENTION

Bonding between boron atoms is more complex than in carbon; for example,both two- and three-center B—B bonds can form. The interaction betweenthese bonding configurations results in as many as 16 bulk allotropes ofboron, composed of icoshedral B₁₂ units, small interstitial clusters,and fused supericosahedra. In contrast, small (n<15) boron clusters formsimple covalent, quasi-planar molecules with carbon-like aromatic oranti-aromatic electronic structure. Recently, it was shown that B40clusters form a cage-like fullerene, further extending the parallelsbetween boron and carbon cluster chemistry.

To date, experimental investigations of nanostructured boron allotropesare notably sparse, partly due to the costly and toxic precursors (e.g.,diborane) typically used. However, numerous theoretical studies haveexamined 2D boron sheets (i.e., borophene). Although these studiespropose various structures, the general class of 2D boron sheets isreferred to as borophene. Based upon a quasi-planar B₇ cluster (FIG.1A), an Aufbau principle has been proposed to construct nanostructuresincluding puckered monolayer sheets (analogous to the relation betweengraphene and the aromatic ring). The stability of these sheets isenhanced by vacancy superstructures or out-of-plane distortions. Earlyreports of multiwall boron nanotubes suggested a layered structure, buttheir atomic-scale structure remains unresolved. It is therefore unknownwhether borophene is experimentally stable and whether the borophenestructure would reflect the simplicity of planar boron clusters or thecomplexity of bulk boron phases.

SUMMARY OF THE INVENTION

In light of the foregoing, it can be an object of the present inventionto provide borophene, related boron allotropes and/or method(s) forpreparation thereof, thereby overcoming various deficiencies andshortcomings of the prior art, including those outlined above. It willbe understood by those skilled in the art that one or more aspects ofthis invention can meet certain objectives, while one or more otheraspects can meet certain other objectives. Each objective may not applyequally, in all its respects, to every aspect of this invention. Assuch, the following objects can be viewed in the alternative withrespect to any one aspect of this invention.

It can be an object of the present invention to provide a method forpreparation of borophene and/or a boron allotrope without resort to useof toxic precursors of the prior art.

It can also be an object of the present invention to provide a method ofusing synthetic conditions to vary and/or control a morphological phaseof such a boron allotrope.

It can also be an object of the present invention, alone or inconjunction with one or more of the preceding objectives, to provide astable, atomically-thin boron allotrope.

Other objects, features, benefits and advantages of the presentinvention will be apparent from this summary and the followingdescriptions of various preferred embodiments, and will be readilyapparent to those skilled in the art having knowledge of elementalboron, nano-dimensioned allotropes thereof and methods for preparation.Such objects, features, benefits and advantages will be apparent fromthe above as taken into conjunction with the accompanying examples,data, figures and all reasonable inferences to be drawn therefrom, aloneor with consideration of the references incorporated herein.

In part, the present invention can be directed to a boron allotropecomprising an elemental boron layer comprising a boron atomic thicknessdimension. In certain non-limiting embodiments, such an allotrope cancomprise a rectangular boron lattice. Alternatively, the presentinvention can be directed to a boron allotrope comprising an elementalboron monolayer comprising a boron atomic thickness dimension, such anallotrope as can be absent a boron compound, a boron alloy or both. Incertain non-limiting embodiments, such an allotrope can comprise arectangular boron lattice. Alternatively, the present invention can bedirected to a metallic boron allotrope comprising an elemental boronlayer comprising a rectangular boron lattice and boron atomic thicknessdimension.

In part, the present invention can also be directed a borophene whichcan comprise an elemental boron layer comprising a boron atomicthickness dimension. In certain non-limiting embodiments, such aborophene can comprise a rectangular boron lattice. In certain suchembodiments, such a borophene can be absent a boron compound, a boronalloy or both. Regardless, such a borophene can be metallic.

In part, the present invention can be directed to an article ofmanufacture comprising a substrate; and a boron allotrope coupled orconnected to, thereon or otherwise contacting such a substrate. Such anallotrope can comprise an elemental boron layer of boron atomscomprising a boron atomic thickness dimension. In certain non-limitingembodiments, such an allotrope can comprise a rectangular boron lattice.As a separate consideration, such a substrate can comprise silver. Incertain such embodiments, such a substrate can comprise single crystalAg(111). As can relate thereto, such a boron allotrope thereon cancomprise at least one of a homogenous boron phase and a striped boronphase. Regardless, such an allotrope can be metallic.

In part, the present invention can also be directed to an articlecomprising a silver substrate; and a metallic boron allotrope thereon.Such an allotrope can comprise an elemental boron monolayer of boronatoms comprising a boron atomic thickness dimension, such an allotropeas can be absent a boron compound, a boron alloy or both. In certainnon-limiting embodiments, such a substrate can comprise single crystalAg(111). In certain such embodiments, such a boron allotrope thereon cancomprise at least one of a homogenous boron phase and a striped boronphase. Regardless, such a substrate and allotrope thereon can be absentoxygen and/or carbon contamination.

In part, the present invention can also be directed to a method ofpreparing an atomically-dimensioned elemental boron allotrope. Such amethod can comprise providing a substrate at a temperature greater thanabout 200° C.; generating elemental boron vapor from a solid elementalboron source; and contacting such a substrate with such boron vapor forat least one of a rate and at a pressure sufficient to deposit, on sucha substrate, a boron allotrope which can comprise an elemental boronlayer comprising a boron atomic thickness dimension. As described belowand illustrated elsewhere herein, such a method can be employed undernegative pressure.

In certain non-limiting embodiments, such a substrate can comprisesilver. In certain such embodiments, such a substrate can comprisesingle crystal Ag(111). As can relate thereto, such a boron allotropethereon can comprise at least one of a homogenous boron phase and astriped boron phase. In certain other non-limiting embodiments, such asubstrate can be at a temperature of about 400° C. to about 750° C. Incertain such embodiments, a higher substrate temperature within such arange can promote such a striped boron phase. In certain othernon-limiting embodiments, as a separate consideration, the rate of boronallotrope deposition can be about 0.01 to about 0.1 layer of elementalboron per minute. In certain such embodiments, a greater deposition ratewithin such a range can promote such a homogenous boron phase.Regardless, such a solid boron source can be about 99.9999% pure. As canrelate thereto, such a method can be employed at a pressure to providedesired boron source purity and/or the absence of carbon and/or oxygenin such a deposited boron allotrope. In certain such embodiments, such amethod can be employed at a pressure less than about 10⁻⁸ mBar. Incertain non-limiting embodiments, such a method can be employed at apressure of about 10⁻⁹ to about 10⁻¹¹ mBar.

In part, the present invention can also be directed to a method ofpreparing an atomically-dimensioned elemental boron allotrope monolayer.Such a method can comprise providing a silver substrate at a temperatureof about 400° C. to about 750° C.; generating elemental boron vapor froma solid elemental boron source; and contacting such a substrate withsuch boron vapor under a pressure at least as low as about 10⁻⁹ mBar,such contact as can be at a rate sufficient to deposit, on such asubstrate, a boron allotrope comprising an elemental boron monolayercomprising a boron atomic thickness dimension, whereby such a methoddoes not provide a boron compound and/or a boron alloy.

In certain non-limiting embodiments, such a substrate can comprisesingle crystal Ag(111). As can relate thereto, such a boron allotropedeposited thereon can comprise at least one of a homogenous boron phaseand a striped boron phase. In certain other non-limiting embodiments,such a substrate temperature can be about 450° C. to about 700° C.; andsuch a rate of deposition can be about 0.01 to about 0.1 layer ofelemental boron per minute. In certain such embodiments, a highersubstrate temperature within such a range can promote such a stripedboron phase. In certain other such embodiments, a greater depositionrate within such a range can promote such a homogenous boron phase.Regardless, such a solid boron source can be about 99.9999% pure; andsuch a deposited boron allotrope can be absent carbon and/or oxygen.

BRIEF DESCRIPTION OF THE DRAWINGS

FIGS. 1A-1L. Growth and atomic scale characterization of borophenesheets. Schematics of (FIG. 1A) distorted B₇ cluster (prior art) and(FIG. 1B) growth setup with atomic structure model and STM topographyrendering. (FIG. 1C) AES spectra of clean Ag(111) before and after borondeposition. (FIGS. 1D-1I) Series of large-scale STM topography (left)and closed-loop dI/dV (right) images of borophene sheets, showing (FIGS.1D, 1E) low coverage (V_(sample)=2.0 V, I_(t)=100 pA), (FIGS. 1F, 1G)medium coverage (V_(sample)=3.5 V, I_(t)=100 pA), and (FIGS. 1H, 1I)high coverage (V_(sample)=3.5 V, I_(t)=100 pA). Regions of homogeneousphase, striped phase island, and striped phase nanoribbon are indicatedwith red, white, and blue arrows, respectively. (FIGS. 1J-1L) STMtopography images showing: (FIG. 1J) Striped phase atomic-scalestructure (V_(sample)=0.1 V, It=1.0 nA). Inset shows rectangular latticewith overlaid lattice vectors. (FIG. 1K) Striped phase with rhombohedral(indicated by white rhombus) and honeycomb (indicated by purple arrow)Moire patterns (V_(sample)=3.5 V, It=100 pA). (FIG. 1L) Striped phaseisland demonstrating carpet-mode growth (V_(sample)=3.5 V, I_(t)=100pA). Inset shows atomic continuity across the Ag(111) step(V_(sample)=−0.5 V, I_(t)=700 pA).

FIGS. 2A-F. Computational prediction of borophene structure andelectronic properties. Top (FIG. 2A) and side (FIG. 2B) views of thelow-energy monolayer structure (unit cell indicated by green box). (FIG.2C) Simulated empty states STM image (V_(sample)=1.0 V), with overlaidatomic structure and unit cell of 0.500 nm by 0.289 nm and (FIG. 2D)experimental STM images (V_(sample)=0.1 V, It=1.0 nA), with overlaidunit cell of 0.51 nm by 0.29 nm. DFT calculated electronic (E) bandstructure (inset: FIG. 2D Brillouin zone) and (FIG. 2F) DOS forfreestanding borophene.

FIGS. 3A-3E. Borophene structural and chemical characterization.Cross-sectional AC-STEM images from (FIG. 3A) HAADF and (FIG. 3B) ABFdetectors. (FIG. 3C) Juxtaposition (left to right) of Si-cappedborophene structure model, simulated ABF image, and magnified ABF image.(FIG. 3D) XPS B is core level spectra and fitted components for sampleswith and without Si capping layers. (FIG. 3E) Angle-resolved XPS dataacquired on Si-capped samples. Inset: schematic showing measurementangle and sample structure determined by angle-resolved XPS.

FIGS. 4A-4B. Scanning tunneling spectroscopy of borophene. (FIG. 4A) STSI-V curves and (FIG. 4B) STS dI/dV spectra (inset: clean Ag(111) dI/dVspectrum) from the borophene sheets, which demonstrate metalliccharacteristics (Feedback loop opened at V_(sample)=1.0 V, I_(t)=1.0nA).

FIGS. 5A-5B. Auger electron spectroscopy analysis and discussion.

(FIG. 5A) AES spectra of clean Ag(111), borophene (−1 ML) grown at 550°C., and a thicker (−4 ML) amorphous boron film. (FIG. 5B) AES spectrumof borophene film with extended range, demonstrating the absence ofimpurity peaks.

FIGS. 6A-6D. Formation of borophene nanoribbons. STM topography imagesof low-flux (−0.02 ML/min), low-coverage deposition at 550° C. showing(FIG. 6A) large scale region with several nanoribbons at step edges(blue arrows) and within terraces (blue arrow). (V_(sample)=1.0 V, h=200pA) (FIG. 6B) Magnified image demonstrating the internal stripedstructure of the nanoribbons. (V_(sample)=0.2 V, I_(t)=500 pA) (FIG. 6C)Atomic scale structure of the nanoribbons, demonstrating theirstructural correspondence with the extended 2D stripe phase sheets.(V_(sample)=−0.1 V, h=5 nA). (FIG. 6D) STS point spectra obtained on therectangular lattice phase regions of the striped boron nanoribbons,demonstrating their metallic characteristics.

FIGS. 7A-7G. Structure analysis and low energy electron diffraction ofborophene. Sequentially acquired STM topography of (FIG. 7A) Ag(111)lattice (V_(sample)=25 mV, I_(t)=1.0 nA) and (B) borophene(V_(sample)=50 mV, I_(t)=2.0 nA). (C) Superimposed fast Fouriertransformations (FFTs) of (FIG. 7A) and (FIG. 7B). (FIG. 7D) FFT oflow-energy monolayer structure model given in FIG. 2 (shown inset). ThisFFT is not scaled to (FIG. 7C), comparison is made relative to theAg(111) spots (circled in red). (FIGS. 7E-7G) LEED patterns acquired at(FIG. 7E) 120 eV, (F) 68.5 eV, and (FIG. 7G) 68.5 eV. (FIG. 7E) and(FIG. 7F) were acquired sequentially on the same 550° C. growth, whereas(FIG. 7G) was acquired on a lower deposition rate growth at 700° C.

FIGS. 8A-8D. Detailed structure and morphology of the homogeneous phase.STM topography images of (FIG. 8A) Atomically resolved homogeneous phasedemonstrating chain morphology, line defect, and Moire structure(V_(sample)=0.1 V, I_(t)=500 pA). (FIG. 8B) Atomically resolved image oftypical chain structure. The short-range Moire pattern is highlightedwith a white rhombus (V_(sample)=0.1 V, I_(t)=3.0 nA). (FIGS. 8C-8D)Atomically resolved images of (FIG. 8C) typical chain structure(V_(sample)=0.1 V, I_(t)=1.0 nA) and (FIG. 8D) magnified region (fromgreen square) demonstrating bias-dependent changes in contrast whichreveal structural relationships between the chains (V_(sample)=−20 mV,I_(t)=1.0 nA).

FIGS. 9A-9F. Detailed structure and morphology of striped phase. (FIG.9A) STM topographic image of the interior of a borophene sheet grown at700° C., demonstrating rhombohedral Moire pattern and regions ofnon-buckled rectangular lattice (V_(sample)=1.2 V, I_(t)=500 pA). (FIG.9B) Atomic resolution image of a borophene striped sheet. Upper inset:FFT of (FIG. 9B). Lower inset: high resolution image (both images atV_(sample)=0.1 V, I_(t)=500 pA). (FIG. 9C) Moire pattern generated fromthe superposition of a strained borophene lattice (4% compressive alongthe a direction) on the Ag(111) plane. The insets show magnified viewsof the Moire trough and crest, detailing the different adsorption sidesin each. (FIG. 9D) Images of 550° C. growth, detailing the interfacialrelationship between a striped phase domain with periodic buckling, anda striped phase domain with nearly complete stripe coverage(V_(sample)=0.2 V, I_(t)=3.0 nA). Purple arrows denote rectangularlattice regions, the blue arrow denotes a non-Moire striped phase, andthe green arrow denotes a Moire striped phase. (FIG. 9E) Atomicresolution image detailing the limited structural coherency and commonstructural motifs between the rectangular and striped phases(V_(sample)=0.1 V, I_(t)=3.0 nA). (FIG. 9F) Atomically resolved image ofstriped phase region generated by the coalescence of multiple islands.The smaller striped domain exhibits a “bullet-like” shapecharacteristics of the homogeneous phase (V_(sample)=0.1 V, I_(t)=1.0nA).

FIG. 10. Computational studies of alloying or boride formation (notobserved). Sample structure model for boron-rich alloy (i.e., matchingthe experimental conditions), showing the restructuring of thenear-surface silver region and segregation of the boron to the surface.

FIGS. 11A-11H. Computational monolayer structures. (FIGS. 11A-11C) top,front, and side views of the atomic structure of a free-standingborophene sheet. Structural relaxation results in decreased corrugationalong the a direction, which results in a decreased unit cell size. Theresulting lattice parameters are a=0.1617 nm (corresponding to ˜⅓ thelattice parameter in FIGS. 2A-2B) and b=0.2865 nm. (FIG. 11D) Phonondispersion of freestanding borophene sheet. The small imaginaryfrequencies near the F point are consistent with instability againstlong-wavelength transversal waves (57). This instability can be fixed bydefects, such as ripples or grain boundaries, which do not allow thesewaves by limiting the size of boron sheets. (FIG. 11E) Calculated strainvs. energy for freestanding borophene sheet. (FIG. 11F) Structure modelcorresponding to a metastable borophene sheet, slightly distorted withrespect to the most stable structure (FIG. 2A). (FIG. 11G) Simulated(V_(sample)=0.1 V, I_(t)=3.0 nA) and (FIG. 11H) experimental(V_(sample)=0.1 V, I_(t)=3.0 nA) STM topography images.

FIGS. 12A-12F. Computational bilayer structures. (FIGS. 12A-12C) Top,side, and frontal views of bilayer borophene structure (not observed).Structural relaxation of the freestanding sheet results in latticeparameters of a=0.490 nm and b=0.288 nm. (FIG. 12D) Calculatedelectronic band structure for freestanding, bilayer borophene and (FIG.12E) associated electronic density of states, both of which showmetallic characteristics. (FIG. 12F) Calculated phonon dispersion forthe bilayer structure. The small region of imaginary frequencies nearthe gamma point is interpreted as in FIG. 11.

FIGS. 13A-13C. Cross-sectional AC-STEM sample prep. SEM images acquiredduring focused ion beam (FIB) milling of a cross-sectional sample forSTEM. Once a suitably flat region is selected, a protective platinumcap/handle is deposited in situ (FIG. 13A). The region surrounding thesample is then milled out using the Ga⁺ ion beam (FIG. 13B). Immediatelyprior to severing the last connection of the sample to the substrate,the sample is welded to the tungsten probe via platinum deposition. Oncethe sample is separated from the substrate, it is transferred via thetungsten probe to the sample post of a semi-circular copper grid. Thesample is welded to the post with platinum deposition, separated fromthe tungsten probe, and successively thinned to <30 nm by the ion beam(FIG. 13C).

FIGS. 14A-14B. Additional AC-STEM data. AC-STEM images acquired on anatomically stepped surface region with (FIG. 14A) HAADF and (FIG. 14B)ABF detectors. The purple arrows are located at the same relativepositions between images, and indicate the presence of light elementalplanar features, which are identified as borophene.

FIGS. 15A-15B. STEM EELS. (A) AC-STEM survey image, with the location ofthe EELS line profile indicated. (FIG. 15B) Plot of EELS spectra fromthe regions indicated in (FIG. 15A). The interface spectrum (red) showsthat a strong boron peak (188 eV) is localized to the area near theinterface, as confirmed by the deconvoluted boron peak (purple). Thesespectra show that the boron is confined to the interface between theAg(111) substrate and the Si capping layer.

FIGS. 16A-16C. Additional ex situ characterization of borophene. (FIG.16A) SEM image of the borophene sheets after ambient exposure. Theisland features observed in STM persist and demonstrate significantcontrast. The small dark features are attributed to boron nanoparticles,and may appear larger than in STM due to charging effects. (FIG. 16B)Ambient peak-force mode AFM image of borophene sheets. Based on theisland morphologies, striped phase (red arrow) and homogeneous phaseislands (blue arrow) are identified. (FIG. 16C) AFM line profilesextracted from (FIG. 16B), showing heights of ˜0.38 nm for the stripedphase island (red arrow) and ˜0.44 nm for the homogeneous phase (bluearrow). These step heights may deviate from theoretical predictions dueto tip-sample interactions and the partial oxidation of the sheets, butare essentially consistent with both theoretical predictions and STEMobservations.

FIGS. 17A-17E. Additional XPS analysis. (FIG. 17A) XPS survey spectrum,demonstrating the presence of Ag, B, Si, C, and 0. (FIG. 17B) Si 2p corelevel data, demonstrating the presence of the almost fully oxidizedamorphous silicon (−103 eV) capping layer adjacent to the Ag 4s peak.(FIG. 17C) C is core level spectrum, demonstrating the absence of anypeaks besides adventitious carbon. (FIG. 17D) Ag 3d core level spectra,showing the silver peaks are unaltered by the B and Si cap deposition.(FIG. 17E) Angle-resolved XPS data showing normalized component peakheights for the species present.

DETAILED DESCRIPTION OF CERTAIN EMBODIMENTS

As can relate to certain non-limiting embodiments of this invention,atomically thin, borophene sheets can be grown under ultrahigh vacuum(UHV) conditions (FIG. 1B) using a solid boron atomic source (e.g.,without limitation, 99.9999% purity) to avoid the difficulties posed bytoxic precursors. An atomically clean Ag(111) substrate can provide awell-defined and inert surface for borophene growth. In situ scanningtunneling microscopy (STM) images show the emergence of planarstructures which can exhibit anisotropic corrugation, consistent withfirst-principles structure prediction. The planar, chemically distinct,and atomically thin nature of such sheets are verified via a suite ofcharacterization techniques. In situ electronic characterizationsupports theoretical predictions that borophene sheets are metallic withhighly anisotropic electronic properties.

As can relate to certain such embodiments, during growth, the substratewas maintained between 450° C. and 700° C. under a boron flux between˜0.01 to ˜0.1 monolayer (ML) per minute. After deposition, in situ Augerelectron spectroscopy (AES; FIG. 1C) revealed a boron KLL peak at thestandard position (180 eV) superimposed on the clean Ag(111) spectrum.No peaks due to contaminants were observed, and none of the distinctivepeak shifts or satellite features characteristic of compound or alloyformation (FIGS. 5A-5B).

After boron deposition at a substrate temperature of 550° C., STMtopography images (FIG. 1D) revealed two distinct boron phases: ahomogeneous phase and a more corrugated “striped” phase (highlightedwith red and white arrows, respectively). Simultaneously acquired dl/dVmaps (where I and V are the tunneling current and voltage, respectively)of the electronic density of states (DOS), given in FIG. 1E, showedstrong electronic contrast between boron sheets and the Ag(111)substrate and increased differentiation between homogeneous and stripedislands. The relative concentration of these phases depends upon thedeposition rate. Low deposition rates favored the striped phase andresulted in the growth of striped phase nanoribbons (blue arrow, alsoFIGS. 6A-6D). At higher deposition rates, more of the homogeneousislands were observed (FIGS. 1F-1G). Increasing growth temperaturesfavored the striped phase, suggesting that the homogeneous phase ismetastable relative to the striped phase. Both phases exhibitedthreefold orientation degeneracy with respect to the substrate, asconfirmed by low-energy electron diffraction (FIGS. 7A-7G). The islandsize for both phases resembles graphene grown on Ag(111). At boroncoverage approaching 1.0 ML, the substrate is completely covered byboron sheets and sparse clusters (FIGS. 1H and 1I).

High-resolution STM images show anisotropic atomic-scale features forboth phases. The homogeneous phase (FIGS. 8A-8D) appears as atomicchains (0.30 nm periodicity) with periodic vertical buckling, ashort-range rhombohedral Moire pattern, and a longer range 1-D Moirepattern (FIGS. 8A-8D). The striped phase (FIG. 1J) consists of arectangular lattice commensurate with regions of striped corrugation.The rectangular structure (inset) is defined by vectors a and b oflengths 0.51 nm (±0.02 nm) and 0.29 nm (±0.02 nm), respectively. Withinthe striped regions, the in-plane periodicity parallel to the a vectoris reduced by the increased out-of-plane corrugation associated with thestripes. However, the periodicities along the stripes match therectangular lattice in the b direction. Further analysis shows that thestriped regions are simple distortions of the rectangular lattice thatmaximize the number of ideal boron adsorption sites (FIGS. 9A-9F). Theformation of these stripes was temperature-dependent, with fewer stripesobserved at 450° C. and almost complete stripe coverage at 700° C. Thisis consistent with a progressive, thermally driven relaxation of therectangular lattice into more favorable adsorption sites.

Rotationally misaligned striped phase islands coalesce via defects thataccommodate the anisotropic corrugations to form a complete monolayer(FIGS. 9A-9F). As shown in FIG. 1K, the striped regions exhibited Moirepatterns with rhombohedral (−8 nm period, marked by white rhombus) or,far less commonly, honeycomb (indicated by purple arrow) symmetry. Theseobservations indicate the possibility of at least two well-definedlong-range structural relationships between borophene and Ag(111). Theborophene superstructure is evidently more complex than planar 2Dmaterials such as BN, which forms a well-defined nanomesh on transitionmetals due to substrate interactions. The mildly attractive B—Aginteractions, result in enhanced corrugation and substrate-stabilizedstructural variation in borophene, providing additional degrees offreedom for functionality beyond conventional 2D materials.

Frequently, borophene growth over the substrate step edges is observed(i.e., “carpet mode” growth, as in FIG. 1L. This continuity of theatomic-scale structure over the step (inset) suggests that the boropheneis structurally distinct from the underlying substrate.

These experimental results are further elucidated by ab initioevolutionary structure prediction with the USPEX algorithm, whichminimizes the thermodynamic potential of the system using densityfunctional theory (DFT). Structures calculated with varyingconcentrations of Ag and B atoms on the Ag(111) substrate show surfacesegregation of B (FIG. 10), indicating that the formation of a B—Agsurface alloy or boride is highly improbable. Additional calculationspredict likely monolayer (FIGS. 11A-11H) and possible bilayer (FIGS.12A-12F) borophene structures on Ag(111), although height measurements,below, supported only a monolayer model.

The lowest energy monolayer structure is shown in FIGS. 2A-2B, and isconstructible from distorted B₇ clusters using the Aufbau principle. Thesymmetry (space group Pmmn) and calculated lattice constants agree wellwith the STM data, with a and b equal to 0.500 nm and 0.289 nm,respectively. Comparison between simulated (FIG. 2C) and experimentalSTM topography images (FIG. 2D, also FIGS. 11A-11H) gives excellentagreement, as does electron diffraction data (FIGS. 7A-7G). Freestandingrelaxation of this structure removes the slight corrugations along the adirection, but preserves the buckling along the b direction (FIGS.11A-11H). The freestanding sheet may exhibit instability againstlong-wavelength transversal thermal vibrations (FIGS. 11A-11H), whichmay contribute to the observed stripe formation and would likely distortthe structure of the borophene sheet upon removal from the growthsubstrate. This substrate-induced stability frames borophene as anintermediate class of templated, covalently bound sheets with propertiesdistinct from conventional 2D materials and more consistent structurethan supported silicon phases.

Electronic band structure calculations (FIG. 2E) within the 2D Brillouinzone of the relaxed, freestanding monolayer (inset) predict metallicconduction (i.e., bands crossing E_(f)) along the F-X and Y-S directions(parallel to the uncorrugated a direction). However, the out-of-planecorrugation along the b direction opens a band gap along the F-Y and S-Xdirections. As a result, borophene is a highly anisotropic metal, whereelectrical conductivity is confined along the chains. The calculated DOS(FIG. 2F) is likewise metallic.

This structure also results in significant mechanical anisotropy (FIGS.11A-11H). Due to the strong, highly coordinated B—B bonds, the in-planeYoung's modulus (a measure of stiffness) is equal to 170 GPa·nm alongthe b direction, and 398 GPa·nm along the a direction, which potentiallyrivals graphene, at 340 GPa·nm. Furthermore, the out-of-plane bucklingresults in negative values for the in-plane Poisson's ratio (equal to−0.04 along a and −0.02 along b), resulting in unusual properties suchas in-plane expansion under tensile strain.

The apparent topographic height of the boron islands in STM dependedupon scanning parameters, with the islands appearing as depressions forsample biases <3.2 V (compare the images in FIGS. 1D-1F). Thisobservation is attributed to the inherent convolution between topographyand electronic structure in STM measurements. Similar inversion isobserved for NaCl islands and graphene on Ag(111). However,cross-sectional, aberration-corrected scanning transmission electronmicroscopy (AC-STEM) unambiguously shows that the boron phase isatomically thin and structurally distinct from the Ag(111) growthsubstrate. AC-STEM sample preparation is detailed in FIGS. 13A-13C.Images acquired with the high-angle annular dark field (HAADF) detector(FIG. 3A) are sensitive to the atomic number Z (contrast ˜Z^(3/2)) andshow minimal contrast at the interface between the Ag(111) substrate andamorphous SiO_(x) capping layer, which is consistent with the lack ofelectron scattering from the low-Z boron. Nevertheless, electron energyloss spectra confirm that the boron lies at the Ag(111) surface (FIGS.15A-15B). Annular bright field (ABF) images (FIGS. 3B and 14A-14B),which are sensitive to light elements such as boron, revealed a planarstructure (indicated by a purple arrow) at this interface. The observedcontrast and structure are consistent with a simulated ABF image of theborophene structure model (FIG. 3C). Measured sheet thicknesses of ˜0.27to ˜0.31 nm match both the monolayer structure model and multiwalledboron nanotubes.

X-ray photoelectron spectroscopy (XPS) measures both sample compositionand the oxidation state of the species present. Although the boropheneislands persisted under ambient conditions (FIGS. 16A-16C), theemergence of higher binding energy features in the XPS B is core levelspectra (FIG. 3D) demonstrate that bare samples (black curve) werepartially oxidized within several hours in ambient. However, thisoxidation was impeded by an amorphous silicon/silicon oxide cappinglayer (red curve) thereon or about, which delayed oxidation for severalweeks (blue curve). The unoxidized, capped and/or encapsulated sample isfit by two Voigt components, which reflect the differences in chemicalenvironment between the low and high buckled atoms. Increasing thephotoelectron detector angle from the sample normal enhances XPS surfacesensitivity, thereby selectively probing the surface and subsurface. Thenormalized, integrated components of angle-resolved XPS spectra onsilicon-capped borophene are plotted in FIG. 3E. With increasingemission angle, the relative intensities of the carbon, silicon, andboron peaks increased, whereas the silver peak diminished. These resultsconfirm the structure shown in the inset schematic, corroborating ourAES, STM, and STEM results. Additional XPS data are given in FIGS.17A-17E.

Examples of the Invention

The following non-limiting examples and data illustrate various aspectsand features relating to borophene, related elemental boron allotropescomprising a layer of boron atoms and/or methods for their preparation,including the assembly of articles comprising borophene and/or such anallotrope, as are available through the synthetic methodologiesdescribed herein. In comparison with the prior art, the presentborophenes, boron allotropes and related articles provide results anddata which are surprising, unexpected and contrary thereto. While theutility of this invention is illustrated through the use of severalarticles, substrates and various borophene/elemental boron allotropelayers which can be used therewith, it will be understood by thoseskilled in the art that comparable results are obtainable with variousother articles, substrates and borophene/boron allotrope layers, as arecommensurate with the scope of this invention.

Materials and Methods.

Materials.

Growth and in situ experiments were conducted in a commercial ultrahighvacuum (UHV) system (base pressure ˜10¹¹ mBar) equipped with apreparation chamber, low-energy electron diffraction (LEED), Augerelectron spectroscopy (AES), and scanning tunneling microscopy (STM).Borophene was grown on single crystal Ag(111) substrates (Mateck,99.999%), which were cleaned via repeated cycles of Ar sputteringfollowed by annealing at 550° C. During growth, samples were heated to˜450° C. to ˜700° C. by a button heater while boron was deposited froman electron beam evaporator (Focus Gmbh.) with a boron rod source (ESPImetals, 99.9999%). The deposition rate was maintained between 0.01 to0.1 ML/min., using dual power sources to deliver 1.8-2.3 kV acceleratingvoltage and 1.5-1.9 A filament current. Using this setup, we obtain ˜50W incident power on a ˜3 mm diameter boron rod. The flux was monitoredusing a faraday cup and periodic calibration on a clean Ag(111) crystalat room temperature. The boron rod was degassed for >6 hours at ˜50 Wheating power, until the chamber pressure was <5.0×10⁻¹⁰ mBar duringoperation. However, depositions at pressures up to ˜2×10⁻⁹ mBar showedno significant alteration in structure or morphology. The siliconcapping layer for ex-situ analysis was deposited from an electron beamevaporator, using a solid silicon rod source (ESPI metals, 99.9999%)using 800 V accelerating voltage and ˜1.8 A filament current to achieve˜15 W power incident on a ˜1.5 mm rod. Degassing occurred at ˜15 W forseveral hours, until the pressure was <5.0×10⁻¹⁰ mBar. The silicon wasdeposited on a cold sample, immediately following retrieval from themicroscope at ˜40 K.

In Situ Characterization.

STM measurements were conducted in an Omicron Nanotechnology VT-SPM ateither 55 K or 300 K. All images presented are at 55 K, unless otherwisenoted. STM measurements were acquired in constant current mode, usingelectrochemically etched W tips degassed in situ. In order to reduceapparent noise in some STM images, a Gaussian filter was applied withradii of 2-3 pixels (i.e., much smaller than the features of interest).Simultaneous dl/dV mapping during scanning and STS spectra were acquiredusing the lock-in technique (10 kHz modulation, 30 mV amplitude). LEEDand in situ AES employed a 4-grid detector (Omicron SpectraLEED), andAES was acquired with incident beam energy of 1.5 keV and a lock-inmodulation of 2.0 V peak-to-peak. AES spectra were charge-corrected tothe position of the elastic primary peak at 1500 eV.

Ex Situ Characterization.

Scanning electron microscopy (SEM) and focused ion beam (FIB) millingwere performed in a FEI Helios system. FIB samples were milled with aGa⁺ ion beam. TEM samples were prepared by FIB and stored under inertatmosphere. Aberration-corrected scanning transmission electronmicroscopy (AC-STEM) imaging and electron energy loss spectroscopy(EELS) were performed using a JEOL ARM 200F microscope equipped with aGatan GIF 863 Tridiem spectrometer system operating at 200 kV. Care wastaken to minimize the electron beam exposure to the sample, in order toreduce electron beam-induced modification of the Ag(111) substrate.High-angle annular dark field (HAADF) and annular bright field (ABF)images were sequentially acquired to provide complementary informationconcerning the sample composition. The Gatan spectrometer was set to anenergy dispersion of 0.5 eV/channel, in order to obtain the best energyresolution at the zero-loss peak. EELS spectra were acquired using thescanning mode to obtain a line profile spectra set to at the same energyas the zero-loss peak, with a 2.0 mm aperture and a 10 ms pixel time per1.0 nm of the sample to increase the signal obtained from the boron (188eV) at the interface. The DeConvEELS software package was applied torectify the EELS spectra by deconvolution with a zero-loss spectrum(thereby acting as a software monochromator). Ambient AFM measurementswere performed on a Bruker Multimode system in Peakforce mode usingsilicon tips with a nominal <10 nm radius of curvature. XPS measurementswere acquired in a Thermo Scientific ESCALAB 250Xi, using an electronflood gun for charge compensation. All XPS data were resolved with ˜0.1eV resolution employing a monochromated Al Ka X-ray source at 1486.7 eV(−400 p.m spot size). General core level spectra were the average of 5scans with a 100-350 ms dwell time, using a pass energy of 15 eV. Forangle-resolved XPS (AR-XPS), the sample stage was tilted with respect tothe x-ray source and photoelectron detector, while single spectra weretaken at each angle. AR-XPS core level spectra data as presented weretaken with pass energies of 50 eV, and the results confirmed with 15 eVpass energy spectra. When using charge compensation, all core levelswere charge corrected to adventitious carbon at 284.8 eV. All subpeakswere fitted with singlets for s photoelectrons, using the software suiteAvantage (Thermo Scientific).

Computational Methods.

Ab initio structure predictions were performed using the USPEX code withsurface prediction module. The Ag(111) substrate was constructed byrectangular and hexagonal lattices, respectively. The lattice vectorsare a=5.004 A and b=2.889 A for rectangular lattice and a=b=5.778 A forhexagonal lattices. The structure searches were conducted with 6, 8, 10,12, 14 and 16 atoms per unit cell for the hexagonal substrate; and 6, 7,8, 10, 12, 14, 16 atoms per cell for the rectangular lattice,respectively. The thickness of the surface layer of two-dimensional (2D)boron and the vacuum were set to 3 Å and 10 Å during the searching, butallowed to change for postprocessing. To investigate the likelihood ofalloy formation, a structure search was performed with the rectangularlattice by varying the number of Ag and B atoms between values of up to8 and 16 atoms. The chemical potential for Ag and B are −2.82 eV/atomand −6.68 eV/atom, respectively. The structural relaxations used theall-electron-projector-augmented wave method as implemented in theVienna ab initio simulation package (VASP). The exchange-correlationenergy was treated within the generalized gradient approximation (GGA),using the functional of Perdew, Burke, and Ernzerhof. The cutoff energyof 450 eV and the uniform F-centered k-points grids with resolution of27c×0.04 A′ were used. Phonon dispersion curves and the simulatedscanning tunneling microscopy (STM) images were computed with thePHONOPY package and Hive code.

The change of energy was calculated in the strain range (between −2% and2% with an increment of 0.5%) to determine the mechanical properties ofthe 2D boron sheet. For a 2D sheet, using the standard Voigt notation,the elastic strain energy per unit area can be expressed asU(ε)=½C ₁₁ε_(xx) ²+½C ₂₂ε_(yy) ² +C ₁₂ε_(xx)ε_(yy)+2C ₆₆ε_(xy) ²where C₁₁, C₂₂, C₁₂, and C₆₆ are components of the elastic modulustensor, corresponding to the second partial derivative of strain energywith respect to strain. The elastic constants can be derived by fittingthe energy curves associated with uniaxial and equibiaxial strains. Whenthis method was applied for penta-graphene, calculated C₁₁ and C_(,2) as288 GPa·nm and −32 GPa·nm were obtained, which are in good agreementwith the reported values (265 and −18 GPa·nm). For the 2D boron sheet,the corresponding C₁₁, C₂₂, C₁₂, and C₆₆ are 398, 170, −7, and 94GPa·nm, respectively (from fitting to FIG. 11E). The in-plane Young'smodulus and Poisson's ratio can be derived from the elastic constants by

${E_{x} = \frac{{C_{11}C_{22}} - {C_{12}C_{21}}}{C_{22}}},{E_{y} = \frac{{C_{11}C_{22}} - {C_{12}C_{21}}}{C_{11}}},{v_{xy} = \frac{C_{21}}{C_{22}}},{v_{yx} = \frac{C_{12}}{C_{11}}}$

Therefore, the Young's modulus and Poisson's ratio are 398 GPa·nm, 0.04in the x direction (i.e., parallel to the a vector), and 170 GPa·nm,−0.02 in the y direction (i.e., parallel to the b vector). Moststrikingly, the in-plane Young's modulus in the x direction iscomparable to or even exceeding that of graphene (340 GPa·nm). Moreover,the 2D boron sheet has a negative C₁₂, leading to be the first 2D boronstructure with negative Poisson's ratio due to the special buckledatomic configuration.

AC-STEM images were simulated with the JEMS software package, usingparameters matching the experimental conditions. An amorphous siliconcapping layer (generated using a-Si Generator) was added to thepredicted structure in order to better model the contrast of the sample.

Example 1

Auger Electron Spectroscopy.

Auger Electron Spectroscopy (AES) in situ verified the cleanliness ofthe surface before deposition, and the chemical identity of the samplesfollowing STM measurements. Auger transitions respond to changes in thelocal chemical environment due to shifts in both the valence and corelevel energies. The formation of compounds generally results in shiftsof at least several eV. Low atomic number elements like boron areespecially susceptible to these shifts.

Representative AES spectra obtained on clean Ag(111), Ag(111) withborophene, and Ag(111) with a thicker (−4-6 ML) amorphous boron film arepresented in FIG. 5A. In the clean Ag(111) spectrum, the characteristicAg MNN spectrum is observed, with sharp peaks at 351 and 356 eV andbroad secondary peaks at 263 and 303 eV. Boron deposition does notnoticeably modify the silver peaks, indicating there was no substantialchemical modification to the surface (i.e., no alloying or compoundformation). Following the deposition of 1 ML at a substrate temperatureof 550° C. (red curve), the boron KLL peak is clearly observed at theexpected 180 eV. Comparison with relatively thick, amorphous boron(deposited at room temperature) shows no extra peaks or peak shifts,which confirms that the measured boron KLL peak position corresponds tounreacted, pure boron.

No additional peaks due to contaminants or oxygen were observed (−500eV, see FIG. 5B), and the absence of carbon was verified through theratio of the silver 266 eV and 304 eV peaks (typical measured value of0.42 before and after boron deposition). The presence of even slightcarbon contamination would significantly distort the Ag peaks in dN/dEspectra. Furthermore, the peak shape does not change followingdeposition, suggesting that no carbon contamination is introduced by theboron source.

Example 2

Low Energy Electron Diffraction.

Spatially averaged atomic structural data is obtained through in situlow-energy electron diffraction (LEED). In LEED, the magnitude of thereciprocal space scattering vector q is given by

$\left| \overset{\rightarrow}{q_{hk}} \right| = \frac{1}{\left| G_{hk} \right|}$where G_(hk) represents the real-space translation vector of the surfaceand h and k represent the indices of the diffraction spot. The sampledregion of reciprocal space is inversely proportional to the electron deBroglie wavelength; therefore, higher electron energy images largerreciprocal lattice vectors (i.e., smaller real space distances).

Reciprocal space structural data are also obtained by performing a fastFourier transform (FFT) on STM images. FIGS. 7A-7B show sequentiallyacquired topography images of the Ag(111) surface and a nearby region ofstriped/rectangular phase. FIG. 7C shows the superposition of FFTs onFIGS. 7A-7B, where the spots related to the Ag(111) and boron structuresare highlighted in red and blue, respectively. The borophene reciprocallattices vectors are indicated. For comparison, FFTs on the calculatedstructure model (FIGS. 2A-2B) are given in FIG. 7D. This calculatedpattern corresponds very well to the experimentally acquired FFT,showing the same relationships between the Ag(111) and borophene spots.Additional spots related to the atomic structure of the borophene areindicated in this model.

These FFT patterns matched the data acquired through LEED, as in FIG.7E, which showed apparent six-fold symmetry for the borophene sheets.The strong diffraction spots circled in red are related to the Ag(111)1×1 surface, whereas the blue spots are related to the structure of theborophene. The apparent six-fold symmetry of these additional spotsresults from the three possible orientation of the boron sheets(observed in STM). The additional spots are related to the Moire patternand buckling of the boron sheets. Lower energy LEED (68.5 eV, FIG. 7F)shows these spots more clearly. The spots highlighted in orange are dueto the a* periodicity, or the formation of the striped domains.Similarly, the spots circled in purple are related to the buckling alongthe b* direction. In the case of the rectangular lattice, thesecorrespond to a 1×3 reconstruction of the relaxed freestanding structureor rectangular 2×(√3/2)R30° with respect to the substrate. The spotscircled in blue correspond to the short-range, rhombohedral Moirepattern in the homogeneous phase, which results in a √3×√3R30° symmetryrelative to the substrate. Many of the lower q-space magnitude (i.e.,larger real-space distance) spots are somewhat less well defined,consistent with the variable periodicities of buckling in the stripephase. The buckling reconstruction parameters for the most orderedstripe patterns in the STM are discussed. Deposition at a slower rate(FIG. 7G) results in a more ordered LEED pattern and elimination of the3×⁻V3R30° spots due to increased conversion of the homogeneous phase tothe striped phase.

Example 3

Structure and Morphology of the Homogeneous Phase.

Several grains of the homogeneous phase are shown in FIGS. 1D-1E.Homogeneous phase islands typically exhibit truncated triangle ortruncated oval (i.e., “bullet-like”) shapes. No homogeneous phaseislands were observed for growth at 700° C., suggesting that hightemperature and/or slow deposition enable the kinetically limited phasetransformation from the homogeneous phase to the striped phase toproceed. Although the homogeneous phase islands often appear featurelessat the large scale, they exhibit multiple scales of Moire pattern,including a ˜0.9 nm rhombohedral pattern, and one-dimensional Moirepatterns with ˜8 nm periodicity, as in FIGS. 8A-8B. The 1D Moire patterncorrelates to variations in the degree of lateral distortion of thechain features in the atomic-scale structure (FIG. 8B). Concurrent withthis chain structure, bias dependent imaging reveals additionalstructure connecting adjacent chains (FIG. 8C), which implies that thebuckled atomic chain structure typically observed likely only capturesthe structure of the highest protruding features. Frequently 1D linedefects are observed, as in FIGS. 8A-8D. These defects sometimes appearas anti-phase boundaries in terms of the inter-chain lateral distortions(i.e., the direction of the lateral distortions is opposite), which mayresult in strain relief or serve to nucleate the proposed phasetransition from homogeneous to striped phase.

Example 4

Striped Phase Structure and Morphology.

Uniaxial stripe structures are common in 2D material heteroepitaxy,resulting from dissimilar symmetries, strain relaxation, andinhomogeneous substrate interactions. Growth at 700° C. results instriped phase islands with a ˜8 nm rhombohedral Moire pattern (FIG. 9A).However, there are always regions with the rectangular lattice withinstriped phase domains, which appear as the dark regions in FIG. 9A. TheMoire pattern continues over these rectangular lattice regions, but isdistorted translationally (shown by the jog in periodicity marked with ablue line) and rotationally (marked by the misalignment of the whitelines). The atomic scale structure within the Moire striped regions isshown in FIG. 9B. In addition to the bucking of adjacent stripe featuresrelative to their neighbors, we observed additional displacement (i.e.,depressions) induced by the large-scale Moire. Similar buckling isassociated with the superposition of domains with dissimilar symmetry.The FFT inset demonstrates the emergence of spots that correspond tothose highlighted in orange and purple in FIG. 7F.

The rhombohedral Moire pattern can be reproduced (FIG. 9C) bysuperimposing a borophene lattice (blue spheres) upon the Ag(111)surface (green spheres), assuming no tilt from the theoreticallycalculated orientation and a 4% in-plane compressive strain along the adirection. The insets show magnified regions of the Moire pattern. Inthe apparent crest (lighter regions, shown in lower inset), we seealignment between the B atoms and the hollow sites, whereas the valleysin the Moire pattern (darker regions, shown in upper inset) showalignment with the Ag atoms. This suggests that the adsorption side ofthe boron atoms drives the corrugation of the striped phase. In BNgrowth on various transition metal substrates, boron is known topreferentially adsorb in the hollow sites of the hexagonal surface,which is consistent with this analysis.

At a growth temperature of 550° C. (FIG. 9D), regions of Moire stripedphase were observed coincident with less regularly striped domains,rotated by 120°. Despite this incongruity in stripe pattern and therotation, the domain boundary remains relatively featureless. Thecongruity between these phases at the atomic scale becomes apparent inFIG. 9E. Regions of the rectangular lattice coincide smoothly with thestriped structure (indicated by white arrow). The rectangular and Moirestriped regions appear to bond readily with one another (indicated bygreen arrow), revealing connected features with three-fold symmetry atthe interface. Nearby, several unit cells of the rectangular lattice areobserved to flow smoothly into the Moire striped phase. The structure ofthe striped phase regions near large gaps between the domains (indicatedby blue arrow) suggests that the striped phase consists of a more highlybuckled lattice. Comparison with the overlaid structure models on therectangular lattice region (upper left) and the striped phase region(lower right) shows that ˜5% strain along the a direction providescoincidence with the striped phase lattice, consistent with the Moirepattern simulated in FIG. 9C. Comparison between the overlaid structuremodel and the striped phase allow us to define a transformation matrix

$\begin{pmatrix}a_{r}^{*} \\b_{r}^{*}\end{pmatrix} = {\begin{pmatrix}{15/2} & {3/2} \\0 & 1\end{pmatrix}\begin{pmatrix}a_{i}^{*} \\b_{i}^{*}\end{pmatrix}}$where the subscripts i and r denote the initial and reconstructedlattice vectors, respectively.

Certain features in STM provide further evidence for conversion from thehomogeneous phase to the striped phase. FIG. 9F shows a region in whicha grain of striped phase is divided into multiple domains. The smallerdomain exhibits an overall shape consistent with that typically found inthe homogeneous phase. Additionally, the shape of the striped phasedomains often closely resembles those of the homogeneous phase (FIGS.1F-1G), but is more compact in size.

Example 5

Scanning Transmission Electron Microscopy/Electron Energy LossSpectroscopy.

The HAADF operation mode works well for materials with different Znumber because the intensity dependence on atomic number is close toZ^(3/2). In FIG. 14A, a HAADF-STEM image shows atomic positions ofsilver. However, light elements such as boron, carbon or oxygen are notdetectable by HAADF imaging. In the JEOL ARM200F microscope bright field(BF) images can be collected at the same time as the HAADF images arerecorded. The BF detector in the microscope has been configured so thatit can register light atoms (extremely weak scattering) using thecircular beam stopper in the center of the optical axis. Under theseconditions, annular bright field (ABF) and HAADF images can be recordedat the same time as well. ABF imaging has proven to be an effective meanof visualization for lighter atoms, such as boron as shown in FIG. 14B.

A STEM-EELS analysis is used to probe the boron signal, the line scan isshown in FIG. 15A. An intensity profile extracted from the EELS spectrumat the B K edge (−188 eV) is correlated with the HAADF image intensityprofile of the same scanned region. During the acquisition, the spectrumof the B K edge intensity is recorded simultaneously with the HAADFsignal. The significant increase in the B K edge intensity profileobserved at the interface between silver/boron/silicon is shown in FIG.15B (red curve, boron deconvolution shown in purple curve). Additionalpeaks between 100 and 150 eV are associated with silicon and siliconoxide, and match those in the capping layer (blue curve).

Example 6

X-Ray Photoelectron Spectroscopy.

Survey spectra on silicon-capped samples (FIG. 17A) show peaks for Ag,B, Si, 0, and adventitious C (which is always present following ambientexposure). Cross-sectional STEM shows an average Si capping layerthickness of 2-4 nm, comparable to the escape depth for photoelectronsin XPS. This allows us to probe the chemistry of the Ag substrate, boronfilm, and Si capping layer simultaneously.

Because pure boron slowly oxidizes under ambient conditions, borophenesamples capped with silicon were compared with ambient-exposed boropheneto determine the oxidation state of the material as deposited, shown inFIG. 3D. The silicon capped sample exhibits one principal feature after2 days, consistent with successful protection of the underlyingborophene. This main peak displays an asymmetric shape composed of twocomponents at 188.2 and 188.8 eV. These are consistent with B—B bondingin pure boron species under two slightly different chemical environments(such as boron atoms interacting directly with the substrate versusthose buckled upwards). When directly exposed to ambient (i.e., no cap),peaks at 187.9, 188.3, and 192.1 eV were observed. This broadening ofthe principal boron peaks, as well as a partial shift towards lowerbinding energies, is typical of boron oxidation, as is the peak at 192.4eV, which likely corresponds to a sub-stoichiometric oxide. Following 3weeks of ambient exposure, the capped sample exhibits similar peaks at187.5, 189, and 192.4 eV. The silicon capping layer was fully oxidized,as shown by the Si 2p peak doublet at ˜103 eV associated with oxideformation on amorphous films (FIG. 17B). C is core level spectra (FIG.17C) showed only adventitious C, indicating that the substrate and Bsource are not contaminated with C (which would otherwise result inadditional peaks for B—C bonds). We observed no splitting of the Ag 3dpeaks (FIG. 17D), which would indicate B—Ag compound formation.

Example 7

Additional angle-resolved XPS data are presented in FIG. 17E, showingthe peak component height versus angle. These data revealed the sametrends observed in FIG. 3E, thus supporting our proposed structuremodel.

As shown above, while all known bulk boron allotropes are semiconductorsat standard conditions. Scanning tunneling spectroscopy (STS) confirmsthe metallic characteristics of borophene through I-V curves (FIG. 4A)and dI/dV spectra (which measure the local electronic DOS, FIG. 4B).These show gapless (i.e., metallic) behavior consistent with thesuperposition between the Ag(111) surface and the predicted filled-statepopulation in borophene (FIG. 2G).

What is claimed is:
 1. A method of preparing an atomically-dimensionedelemental boron allotrope, said method comprising: providing a substrateat a temperature greater than about 200° C.; generating elemental boronvapor from a solid elemental boron source; and contacting said substratewith said boron vapor for at least one of a rate and at a pressuresufficient to deposit on said substrate a boron allotrope comprising anelemental boron layer comprising a boron atomic thickness dimension,said method under negative pressure.
 2. The method of claim 1 whereinsaid substrate comprises Ag.
 3. The method of claim 2 wherein saidsubstrate comprises single crystal Ag(111).
 4. The method of claim 3wherein said allotrope comprises at least one of a homogeneous boronphase and a striped boron phase.
 5. The method of claim 4 wherein saidsubstrate temperature is about 400° C. to about 750° C.
 6. The method ofclaim 5 wherein a higher substrate temperature within said temperaturerange promotes said striped boron phase.
 7. The method of claim 1wherein said rate of deposition is about 0.01 to about 0.1 layer ofelemental boron per minute.
 8. The method of claim 7 wherein a greaterdeposition rate within said deposition rate range promotes saidhomogeneous boron phase.
 9. The method of claim 1 wherein said solidboron source is about 99.9999% pure.
 10. The method of claim 9 whereinsaid deposited boron allotrope is absent carbon and oxygen.
 11. Themethod of claim 1 at a pressure of about 10⁻⁹ to about 10⁻¹¹ mBar.
 12. Amethod of preparing an atomically-dimensioned elemental boron allotropemonolayer, said method comprising: providing a silver substrate at atemperature of about 400° C. to about 750° C.; generating elementalboron vapor from a solid elemental boron source; and contacting saidsubstrate with said boron vapor under a pressure at least as low asabout 10⁻⁹ mBar, said contact a rate sufficient to deposit on saidsubstrate a boron allotrope comprising an elemental boron monolayercomprising a boron atomic thickness dimension, whereby said method doesnot provide a boron compound or a boron alloy.
 13. The method of claim12 wherein said substrate comprises single crystal Ag(111).
 14. Themethod of claim 13 wherein said allotrope comprises at least one of ahomogeneous boron phase and a striped boron phase.
 15. The method ofclaim 14 wherein said substrate temperature is about 450° C. to about700° C.; and said rate of deposition is about 0.01 to about 0.1 layer ofelemental boron per minute.
 16. The method of claim 15 wherein a highersubstrate temperature within said temperature range promotes saidstriped boron phase.
 17. The method of claim 15 wherein a greaterdeposition rate within said deposition rate range promotes saidhomogeneous boron phase.
 18. The method of claim 12 wherein said solidboron source is about 99.9999% pure, and said deposited boron allotropeis absent carbon and oxygen.